Zn-Al Alloy Having Excellent High-Speed Deformation Properties and Process For Producing the Same

ABSTRACT

A Zn—Al alloy excellent in static deformability as well as dynamic deformability and applicable to large-sized structures, and a method for production thereof. The alloy contains 30-99% Zn, with the remainder being Al and inevitable impurities, and has a metallographic structure in which the α phase or α′ phase having an average grain size no larger than 5 μm contains the β phase finely dispersed therein, the Al inclusions have a maximum equivalent circle diameter no larger than 50 μm and are free of pores no smaller than 0.5 mm in terms of equivalent circle diameter, and the macrosegregation of Al is less than 3.0% and the microsegregation of Al is less than 2.0%. (% means mass %.)

TECHNICAL FIELD

The present invention relates to a Zn—Al alloy and a method for production thereof. More particularly, the present invention relates to a Zn—Al alloy which exhibits excellent deforming capability when it experiences rapidly applied stresses, and also to a method for production thereof. The Zn—Al alloy disclosed herein is useful for the seismic isolator or damper which permits any building to follow its vibration and strain caused by winds and earthquakes.

BACKGROUND ART

There have been contrived several seismic isolators or dampers which absorb or follow strains due to winds or earthquakes. They include lead dampers, vibration proof rubber, oil dampers, and vibration controlling steel sheets, for example, LYP (extra low yield point steel). Unfortunately, the vibration proof rubber, which deteriorates with time, is not suitable for use as the seismic isolator or damper of buildings which are required to have long-term durability. Oil dampers, which need periodic maintenance, are not suitable for use as the seismic isolator or damper of buildings, like the vibration proof rubber. The vibration controlling steel sheets, for example, LYP undergo work hardening due to permanent set and deteriorate due to repeated loads, with the result that not only do they decrease in energy absorbing capacity but they also become too hard to isolate vibration propagating to structures. Thus they are limited in use for the seismic isolators or dampers.

By contrast, the lead damper shown in FIG. 1 is soft enough to follow slow vibrations of 0.1 to 10 Hz induced by earthquakes or winds and are hardly subject to deterioration due to expansion and contraction. It is now in common use as a seismic isolator or damper for buildings. Incidentally, FIG. 1 shows a lead casting (1), a homogen welded part (2), and a steel sheet (3).

Such large dampers, however, are heavy and unhandy. Moreover, they need a special technique for their connection to the structure or members attached thereto on account of the low yield point of lead, which is about 5 MPa. In addition to these disadvantages, lead toxicity restricts their use in the field of building.

Under circumstances mentioned above, there has been a demand for a new vibration damping metal which is non-toxic and suitable for small light-weight devices. A notable metal as a substitute for lead is the Zn—Al alloy which exhibits superplasticity.

A recent report about the Zn—Al alloy mentioned above says that it exhibits such superplasticity as to follow a strain rate of 1×10⁻⁴ s⁻¹ at 373K or about 100° C. if it contains 22% Al and has a fine crystalline structure of the order of nanometers. (See Non-Patent Document 1.) This alloy, however, cannot actually be applied to the seismic isolator for buildings which needs elongation at room temperature.

A report about Zn-22% Al-2% Cu alloy says that it exhibits superplasticity at room temperature if it undergoes isothermal water cooling and subsequent cold working to yield a structure composed of α-phase and β-phase precipitating therein. (See Non-Patent Document 2.)

The report also reveals that the alloy marked an elongation of 135% or 160% at the maximum. However, it mentions nothing about whether or not it exhibits such a large elongation at room temperature after warm working. The alloy should exhibit a much larger elongation (say, 180% or above) at room temperature even after cold working if it is to have seismic isolating or damping performance which is good enough to supersede lead dampers.

There is a report about an experiment with a small sample of Zn-22% Al alloy that exhibits superplasticity at room temperature. (See Non-Patent Document 3.) The report specifically reveals that a sample of Zn-22% Al alloy in cylindrical form having an initial grain size of 1 to 15 μm in its metallographic structure changes to have the final structure with the minimum grain size of 0.1 to 0.5 μm at the center after cold deformation or strong twisting under a high pressure (5 GPa).

However, the sample mentioned in the report lacks cross-sectional uniformity in structure. It has a fine structure with superplasticity at its center but it has a coarse granular structure without superplasticity in its periphery. In addition, the intense twisting deformation can only be applied to a very small sample, say 15 mm in diameter and 0.3 mm in thickness. Thus, it is difficult to turn any material, such as seismic isolator which is subject to large loads, into one which has the above-mentioned fine structure throughout it. Consequently, it is impossible to obtain any material which entirely exhibits superplasticity.

The present inventors have been carrying out investigation on Zn—Al alloys for their improvement in characteristic properties. And, they have proposed, as a part of the results of their investigation, a new Zn—Al alloy suitable for use as a building material on account of its uniform, stable ultrafine structure and its ability to exhibit considerable elongation resembling superplasticity even at room temperature. (See Patent Document 1.) The Zn—Al alloy of practical size resulting from the development of the new technology still has problems to be solved despite its superplasticity at room temperature. It is good in deformability at a low strain rate of about 10⁻³ s⁻¹ (referred to as “static deformation” hereinafter) and hence exhibits good superplasticity at room temperature but it is poor in deformability at a comparatively high strain rate of about 10⁻¹ s⁻¹ (referred to as “dynamic deformability” hereinafter). This property becomes significant as the ingot size increases.

Patent Document 1:

Japanese Patent Laid-open No. Hei-11-222643 (claims)

Non-Patent Document 1:

R. S. Mishra et al., The observation of tensile superplasticity in nanocrystalline materials; Nanostruct Mater. vol. 9, No. 1/8, p. 473-476 (1997)

Non-Patent Document 2:

G. Toress-Viallsenor et al., A reinvestigation of the mechanical history on superplasticity of Zn-22Al-2Cu at room temperature, Material. Science. Forum vol., 243/245 p. 553 (1997) Non-Patent Document 3: M Furukawa et al., Fabrication of submicrometer-grained Zn-22% Al by torsion straining; J. Mater. Res. vol. 11, No. 9, p 2128 (1996)

DISCLOSURE OF THE INVENTION Problems to be Solved by the Invention

The present invention was completed in view of the foregoing. It is an object of the present invention to provide a new Zn—Al alloy and a method for production thereof, said alloy being superior in static deformability as well as dynamic deformability and applicable to large structures.

Means for Solving the Problems

The present inventors carried out their investigations from the viewpoint of improving the deformation properties at a high strain rate of 10⁻¹ s⁻¹. The result of investigations revealed that the object is achieved by controlling the ultrafine structure as disclosed in Patent Document 1, and also by reducing coarse Al inclusions and macro- and micro-segregation as disclosed in Japanese Patent Application No. 2002-328106.

After their further investigation, the present inventors found a new Zn—Al alloy (and a method for production thereof) which produces the same effect as disclosed before even in the absence of strict control for the ultrafine structure. This finding shows that the Zn—Al alloy exhibits excellent dynamic deformability if it is produced in such a way that its structure does not contain inevitable pores no smaller than 0.5 mm, even though the average grain size in β phase is larger than 0.05 μm. The foregoing is contrary to the conventional belief that β phase dispersed in α phase or α′ phase should have a grain size no larger than 0.05 μm if the alloy is to exhibit superplasticity. This finding has led to the concept of the present invention.

The Zn—Al alloy according to the present invention is one which contains 30-99% Zn, with the remainder being Al and inevitable impurities. (“%” means “mass %” hereinafter.) It has a metallographic structure in which α phase (or α′ phase) contains finely dispersed β phase and has an average grain size no larger than 5 μm, Al inclusions have a maximum size no larger than 50 μm (in terms of equivalent circle diameter), there exist no pores no smaller than 0.5 mm (in terms of equivalent circle diameter), macrosegregation of Al is less than 3.0, and microsegregation of Al is less than 2.0%.

The present invention does not elucidate the reason why the alloy free of pores in specific size mentioned above exhibits excellent dynamic deformability even though the average grain size in the β phase is slightly large. It is considered that fracture starts at pores and hence reduction in pore size leads to an increase in toughness.

FIG. 5 shows the metallographic structure composed of α phase and β phase finely dispersed therein. The latter, which is finely dispersed in the former, is usually composed of some crystal grains, as shown in FIG. 6. Since any alloy is poor in dynamic deformability if it contains β phase having a large average grain size, the average grain size of β phase should preferably be smaller than 3 μm, more preferably smaller than 0.1 μm, particularly smaller than 0.05 μm. Nevertheless, it has been found that excellent dynamic deformability can be achieved even though the average grain size is larger than 3 μm, as mentioned later.

However, the average grain size should preferably be smaller than 5 μm; otherwise, the finely dispersed β phase becomes coarse. Incidentally, the β phase may be composed of only one crystal grain or a plurality of crystal grains. The preferable grain size is smaller than 5 μm in the former case and smaller than 10 μm in the latter case.

The present invention also covers a method of producing the Zn—Al alloy excellent in rapid deformation properties. The method comprises:

a step of casting a molten Zn—Al alloy into a mold while isolating the melt from the ambient atmosphere, a step of cooling the mold after casting at an average cooling rate no lower than 0.25° C./s in the temperature range from 425 to 375° C. and no lower than 0.020° C./s in the temperature range from 275 to 250° C., a step of reheating which includes keeping hot at 350° C. or above and subsequent quenching, a step of blooming at 275° C. or below, and a step of warm working at 275° C. or below.

EFFECT OF THE INVENTION

The Zn—Al alloy according to the present invention is excellent in static deformability as well as dynamic deformability, and hence it is suitable for use as a seismic isolating material of large structures.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic diagram showing the construction of a conventional lead damper.

FIG. 2 is a graph showing the relation between the temperature in a heating furnace and the temperature of a bloom, which is observed when a 150-kg ingot undergoes atmospheric heating.

FIG. 3 is a phase diagram of the Zn—Al alloy.

FIG. 4 is a graph showing an example of the cooling curve (representing the change with time in temperature inside the ingot).

FIG. 5 is an example of SEM photograph (×5000) showing the metallographic structure (α phase containing β phase dispersed therein) according to the present invention.

FIG. 6 is an example of enlarged photograph showing the β phase of the metallographic structure (α phase containing β phase dispersed therein) according to the present invention.

FIG. 7 is a photograph which schematically shows the positions for measurement of the average grain size of α phase in the metallographic structure according to the present invention.

EXPLANATION OF SYMBOLS

-   1 . . . . Lead casting -   2 . . . . Homogen welded part -   3 . . . . Steel sheet

BEST MODE FOR CARRYING OUT THE INVENTION

The Zn—Al alloy of the present invention should have a specific structure as follows. For the Zn—Al alloy to exhibit superplasticity, it should have a structure such that the β phase is dispersed and precipitated in the α phase or α′ phase. [This structure will be referred to as “β-dispersed α phase” hereinafter.] The β-dispersed α phase is entirely different from α phase free of β phase precipitating therein, and it exhibits elongation in excess of 200% due to plastic deformation resulting from movement of crystal grains.

The Zn—Al alloys vary in metallographic structure depending on the Zn content even though they meet the foregoing requirements. The Zn—Al alloy containing 30-80 wt % Zn, with the remainder being Al and inevitable impurities, has a single macroscopic structure composed of α phase or α′ phase containing β phase finely dispersed therein.

By contrast, the Zn—Al alloy with a Zn content exceeding 80 wt % inevitably has a mixed structure composed of α phase and β phase as is evident from FIG. 3. The Zn—Al alloy containing 80-99 wt % Zn, which is specified in the present invention, has a two-phase mixed structure composed of α phase (or α′ phase) and coarse β phase finely dispersed therein which has a grain size of about 10 μm.

The Zn—Al alloy containing 80-99 wt % Zn, with the remainder being Al and inevitable impurities, exhibits an elongation in excess of 160% as a whole. This is because the β-dispersed α phase exhibits an elongation in excess of 200%, thereby avoiding stress concentration at the grain boundary of the β phase, even though the coarse β phase simply exhibits a ductility of about 65% (which is due to recovery at normal temperature).

By contrast, an alloy of two-phase structure composed of α phase (without β precipitation therein) and β phase does not exhibit superplasticity although α phase and β phase individually exhibit ductility. In addition, the coarse β phase becomes stable in deformation resistance on account of recovery of dislocation at normal temperature but it has a low elongation of about 65%; therefore, the two-phase structure composed of α phase (without β precipitation therein) and β phase merely exhibits a low elongation of about 68% as a whole.

The absence of α phase (free of β phase precipitating therein) and coarse β phase is desirable; however, the presence of β phase (to an extent not harmful to the effect of the present invention) is permissible in the β-dispersed α phase which exhibits superplasticity. Here, the coarse β phase contains the lamella structure mentioned later.

For the Zn—Al alloy of the present invention to exhibit superplasticity (static deformability) in excess of 160% at room temperature, it should have a metallographic structure in which the α phase (or α′ phase) contains the β phase finely dispersed therein. The average grain size in the α phase or α′ phase should preferably be no larger than 5 μm, particularly no larger than 3.5 μm, because the smaller grain size readily induces superplasticity.

For the Zn—Al alloy to have good dynamic deformability, it is necessary to reduce the amount of coarse Al inclusions and macro- and micro-segregations and to reduce the pore size.

<Maximum Diameter of Al Inclusions: No Larger than 50 μm in Equivalent Circle Diameter>

Coarse Al inclusions start off fracture, thereby adversely affecting dynamic deformability as well as static deformability. Therefore, the amount of coarse Al inclusions should be minimal. The present invention specifies that Al inclusions should not be larger than 50 μm (preferably 20 μm) in equivalent circle diameter. Incidentally, Al inclusions are composed mainly of Al₂O₃.

<Macrosegregation of Al: Less than 3.0%, Microsegregation of Al: Less than 2.0%>

Macrosegregation is defined as segregation that takes place throughout an ingot. The present invention requires that there should be no difference larger than 3.0% (preferably 2.0%) between the Al concentration in macrosegregation and the average Al concentration in the ingot. (That is, the amount of macrosegregation should be less than 3.0%.) On the other hand, microsegregation is defined as segregation that takes place in a very small portion (a few micrometers) consisting of several crystal grains. The present invention requires that there should be no difference larger than 2.0% (preferably 1.0%) between the Al concentration in microsegregation and the average Al concentration in the ingot.

<Pore Size>

Like coarse Al inclusions, coarse pores also start off fracture, thereby adversely affecting dynamic deformability as well as static deformability. Therefore, the amount of coarse pores should be minimal. The present invention specifies that pores should not be larger than 0.5 mm (preferably 0.3 mm) in equivalent circle diameter. The maximum pore diameter should desirably be smaller than 0.3 mm if the β phase has a large average grain size.

Meeting the above-mentioned requirements for macrosegregation and microsegregation and the minimum existence of pores is essential for good dynamic deformability (or rapid deformation). Without any one of them, the object of the present invention will not be achieved.

The Zn—Al alloy of the present invention should have a chemical composition as follows. The content of Zn is 30 to 99%, preferably 50 to 99%, more preferably 70 to 99%, and the remainder consists of Al and inevitable impurities. Preferable among them is a Zn-22% Al eutectic alloy. It permits easy structure control and exhibits superplasticity readily because it has the eutectic point at the Al content of 22%, as shown in FIG. 3 (which is a phase diagram of the Zn—Al alloy).

The Zn—Al alloy tends to decrease in elongation according as the Zn content decreases in the above-mentioned range even though the amount of β segregation decreases and plastic deformation takes place due to movement of crystal grains. With a Zn content less than 30%, it will not produce elongation exceeding 100% even if it is treated under the conditions specified in the present invention. Incidentally, FIG. 3 shows the α phase, which is a region composed mainly of Al crystals of face centered cubic lattice structure, the α′ phase, which is a region composed mainly of Zn crystals of face centered cubic lattice structure, the β phase, which is a region composed mainly of Zn crystals of hexagonal close-packed lattice structure, and the liquid represented by L.

Meeting the above-mentioned requirements, the Zn—Al alloy of the present invention may be incorporated with reinforcing elements, such as Cu, Si, Mn, and Mg, in an amount harmless to the hysteresis stability, so that it keeps its stationary stress almost unchanged by the amount of working and the rate of strain. For improvement in elongation, it may be additionally incorporated with Zr or TiB which makes crystals finer.

The Zn—Al alloy meeting the above-mentioned requirements may be obtained efficiently in the following manner.

<Casting Step: the Melt of Zn—Al Alloy should be Kept Isolated from the Ambient Atmosphere During its Casting into a Mold.>

Casting in this manner excludes oxygen which otherwise binds to the alloy melt, thereby preventing Al₂O₃ from becoming coarser. This keeps Al inclusions smaller than 50 μm in equivalent circle diameter at the largest. Isolation from the ambient atmosphere may be effectively accomplished by casting in a vacuum or argon gas or by immersing a gas injecting nozzle into the melt.

<Mold Cooling Step (I) after Casting: with an Average Cooling Rate No Lower than 0.25° C./s in the Temperature Range of 425 to 375° C.>

The mold cooling step subsequent to casting needs an average cooling rate no lower than 0.25° C./s while the mold is cooling from 425 to 375° C. This temperature range corresponds to the solid-liquid dual phase region. Cooling in this manner suppresses the coarse solid structure that causes macrosegregation. In other words, the comparatively rapid cooling in the specified temperature range gives rise to coarse Al precipitates, which prevent the formation of coarse solid structure. The average cooling rate should preferably be no lower than 0.30° C./s.

<Mold Cooling Step (II) after Casting: with an Average Cooling Rate No Lower than 0.020° C./s in the Temperature Range of 275 to 250° C.>

The mold cooling step subsequent to casting needs an average cooling rate no lower than 0.020° C./s while the mold is cooling from 275 to 250° C. This temperature range corresponds to the α+β dual phase region. Cooling in this manner prevents the precipitates of the β phase from becoming coarser and suppresses microsegregation resulting mainly from the coarse β phase in the α phase. In other words, the comparatively rapid cooling in the specified temperature range prevents the precipitates of Zn and Al from becoming coarser, thereby suppressing the formation of the coarse β phase (larger than 10 μm) and giving rise to the finely dispersed β phase. The average cooling rate should preferably be no lower than 0.025° C./s.

<Reheating Step: Heating and Keeping at 350° C. or Above and Subsequent Quenching>

Although rapid cooling in the mold suppresses the formation of coarse solid structure to some extent as mentioned above, the cooling step should be followed by reheating for homogenization to enhance its effect.

For sufficient homogenization, the soaking temperature should be no lower than 350° C. but lower than 390° C. The ingot will melt at a temperature above 390° C.

About one hour of reheating at the temperature specified above will be enough to homogenize small ingots weighing 50 kg or less. However, a much longer time will be necessary to heat large ingots weighing 150 kg or more entirely to 350° C. or above.

FIG. 2 is a graph showing the relation between the furnace temperature and the temperature which was observed during atmospheric heating of a 150-kg ingot of Zn—Al alloy. It suggests that heating for 8 hours is necessary to raise the temperature to 350° C. or above. This is because the ingot absorbs a large amount of heat from outside when the β phase particles dissolve again in the α phase matrix. Therefore, a large ingot inevitably needs a prolonged atmospheric heating. A possible solution to this problem is high-frequency induction heating in place of atmospheric heating. However, induction heating of large ingots leads to a higher production cost.

Reheating at 350° C. or above for a prescribed period of time is followed by quenching. This quenching is necessary for the alloy to have a specific metallographic structure free of microsegregation, with the β phase being confined in the α phase. Quenching should lead to room temperature or blooming temperature (mentioned later).

Quenching of the alloy with the above-mentioned metallographic structure prevents transfer of the α′ phase to the stable α phase, thereby preventing diffusion of the β phase to such an extent that the alloy separates into two phases in a macro level. The result is that the β phase remains precipitated in the α phase. Thus there is obtained the β-dispersed α phase which exhibits superplasticity. Incidentally, “quenching” mentioned above means cooling (preferably water cooling) at a cooling rate no lower than 10° C./s. Furnace cooling (slower than 0.1° C./s) or air cooling (slower than 10° C./s) is not desirable because it causes the diffusion of β phase that leads to lamellar structure. The overall lamellar structure that occurs in this stage prevents the α phase or β phase from becoming fine sufficiently in the subsequent working (mentioned later), especially the one with a low working ratio. The resulting alloy will merely have a limited elongation of 100 to 140%; it does not exhibit a high elongation exceeding 160%.

Although the lamella structure should preferably be absent, it may be partly present in the entire structure. In this case, it should preferably have a size of about 30 μm and an areal ratio of about 20% or less.

<Blooming Step: To be Performed at 275° C. or Below>

In the stage after soaking and subsequent quenching, there exists the β-diffused α phase, in which the α phase or α′ phase has a grain size of about 10 to 2 μm and the β phase has a grain size smaller than 5 μm, particularly about 0.05 to 0.1 μm. The structure like this exhibits an elongation higher than 180% close to superplasticity at high temperatures from about 100 to 150° C. However, this does not hold true at room temperature.

For the alloy to exhibit a high elongation close to superplasticity at room temperature, soaking and quenching should be followed by application of external physical force which will reduce the grain size in the α or α′ phase or in the β phase existing in the α or α′ phase and also crush pores. To this end, the steps of reheating and quenching (mentioned above) should always be followed by blooming (or forging) at 275° C. or below and subsequent warm working.

Blooming temperature should not exceed 275° C.; otherwise, transformation takes place to separate the β-dispersed α phase into two-phase structure composed of α or α′ phase and β phase, as shown in FIG. 3. Blooming temperature should preferably be lower than 200° C. but higher than 100° C. to avoid cracking.

Incidentally, the blooming and warm working mentioned above sufficiently reduce the grain size and hence eliminate the necessity of additional cold working.

According to the present invention, blooming may be followed by warm working immediately at a high temperature without cooling or after cooling to room temperature. In the latter case, the cooling rate should be no lower than about 3° C./s so that the resulting β-dispersed α phase is fixed in the same way as in cooling that follows reheating. Water cooling is desirable for such a cooling rate.

<Warm Working Step: To be Performed at 275° C. or Below>

The warm working should be performed at 275° C. or below. This is because warm working at high temperatures exceeding 275° C. brings about transformation in structure and causes the resulting β-dispersed α phase to separate into two-phase structure of α or α′ phase and β phase. A desirable temperature for warm working is no higher than 200° C. but no lower than 100° C. because warm working at an excessively low temperature will cause cracking.

The warm working mentioned above is not specifically restricted so long as it applies external force large enough to reduce the grain size. It includes, for example, forging, extrusion, and drawing.

The warm working should be followed by cooling (desirably water cooling) to room temperature at a rate of about 3° C./s. This cooling fixes the β-dispersed α phase in the same way as in cooling that follows reheating. Excessively slow cooling makes the β-dispersed α phase coarse, so that the resulting alloy will not exhibit superplasticity at room temperature.

Since the Zn—Al alloy of the present invention is as hard as or slightly softer than mild steel, it can be readily joined to building structures by bolting or riveting. However, in case of soldering or the like which involves heating, it should be kept below 250° C., preferably below 100° C. This is because heating above 250° C. causes the structure to undergo transformation and heating above 100° C. without subsequent quenching makes the valuable fine structure coarse, and the resulting alloy will not exhibit an elongation exceeding 160% at room temperature, as mentioned above.

The invention will be described in more detail with reference to the following examples, which are not intended to restrict the scope thereof. Any modification and change may be made to the examples without altering the scope of the invention.

EXAMPLE 1 Preparation of Zn—Al Alloy

Various ingots (each weighing 180 kg) of Zn-22% Al alloy (with less than 0.5% impurities in total) liable to macrosegregation were prepared by casting into an air-cooled or water-cooled iron or copper mold measuring 200×350 mm in cross section, except that sample No. 12 (shown in Table 1) was prepared by continuous casting into a water-cooled copper mold measuring 200×200 mm in cross section.

The cooling behavior of each ingot was observed with a thermocouple placed in each ingot at the center of the cross section 300 mm above the bottom. FIG. 4 shows cooling curves representing the temperature change with time in the ingot. From the cooling curves were calculated the average cooling rate 1 in the solid-liquid dual phase region (425-375° C.) and the average cooling rate 2 at the temperature (275-250° C.) at which the β phase begins to precipitate. Incidentally, the melt was isolated from the ambient atmosphere during casting by keeping the mold inside and spout sealed with argon gas. In the case of continuous casting, sealing was achieved by immersing the nozzle in the melt.

The resulting ingot of Zn—Al alloy was reheated (for soaking) and kept for 1 hour or 8 hours in an atmospheric furnace at the temperature shown in Table 1. The duration of reheating is a length of time after the slab temperature had reached a prescribed temperature which was measured with a thermocouple in contact with the slab surface in the atmospheric furnace.

Immediately after removal from the furnace, the reheated ingot was cooled with water down to the blooming temperature. The cooled ingot was bloomed (forged) into a block measuring 350×200×450 mm at the temperature shown in Table 1 by means of a 400-ton hydraulic press. This step was followed by water cooling. (The ingot without blooming was water-cooled to room temperature after reheating.) Finally, the block underwent warm working (isothermal rolling) and subsequent water cooling. Thus there was obtained a 20-mm thick alloy plate.

TABLE 1 Average Average Reheating Warm Cooling rate Amount cooling cooling condition Blooming working after of melt Sealing rate 1 rate 2 Temp. Duration (forging) (rolling) warm working Sample No. (kg/ch) Method of casting method (° C./s) (° C./s) (° C.) (h) temp. (° C.) temp. (° C.) (° C./s) 1 180 Air-cooled iron mold Argon 0.160 0.010 350 8 250 200 0.1 2 180 Water-cooled iron mold Argon 0.200 0.017 350 8 250 200 10 3 180 Water-cooled copper mold Argon 0.300 0.025 350 8 None 200 10 4 180 Water-cooled copper mold Argon 0.300 0.025 300 8 None 200 10 5 180 Water-cooled copper mold Argon 0.300 0.025 350 8 250 250 10 6 180 Water-cooled copper mold Argon 0.300 0.025 350 8 250 350 10 7 180 Water-cooled copper mold Argon 0.300 0.025 350 8 150 250 10 8 180 Water-cooled copper mold Argon 0.300 0.025 350 8 150 250 10 9 180 Water-cooled copper mold Argon 0.300 0.025 350 8 350 250 10 10 180 Water-cooled copper mold Argon 0.190 0.010 350 8 250 250 10 11 180 Water-cooled copper mold None 0.300 0.025 350 8 250 250 10 12 180 Water-cooled copper mold Nozzle 0.500 0.033 350 8 250 250 10 immersion 13 180 Water-cooled copper mold Argon 0.300 0.025 300 8 250 250 10 14 180 Water-cooled copper mold Nitrogen 0.300 0.025 350 8 250 250 10 15 180 Water-cooled copper mold Argon 0.300 0.025 350 1 250 250 10 16 180 Water-cooled copper mold None 0.300 0.025 350 8 250 250 10 17 180 Water-cooled copper mold Argon 0.300 0.025 350 8 250 270 10

The thus obtained samples of Zn—Al alloy were tested for characteristic properties in the following manner.

[Evaluation of Characteristic Properties]

Each sample of the alloy plates obtained as mentioned above was examined under an electron microscope to measure the grain size of α phase (including α′ phase).

To be more specific, each alloy sample, with its surface buffed and etched, was observed under an SEM (scanning electron microscope) with a magnification of ×5000. Three fields of view were photographed for each sample. Each photograph was scribed with arbitrary three straight lines (100 μm long), and the length of the straight line passing through each α phase was regarded as the grain size of the α phase. An average of three measurements was obtained as an equivalent circle diameter. The same procedure as above was repeated differently according to the size of α phase by using the magnification of ×1000 to ×10000 and scribing straight lines with a length of 500 μm or 50 μm. (See FIG. 7.)

A specimen (10 mm thick, conforming to JIS No. 5) was taken from each sample. The specimen underwent tensile test under the following conditions.

Gauge length: 50 mm Cross head speed: 5 mm/min (corresponding to static deformability at a strain rate of 1.67×10⁻³ s⁻¹) Cross head speed: 250 mm/min (corresponding to dynamic deformability at a strain rate of 8.33×10⁻² s⁻¹) Tensile strength (TS) and elongation (El) at break were measured. In this way, each sample was examined for static characteristics (TS and El in slow deformation) and dynamic characteristics (TS and El in rapid deformation).

The maximum grain size of Al inclusions was measured as follows. A sample of rolled slab (18 mm thick), with its surface buffed and polished, was observed under an optical microscope with a magnification of ×1000. The point of observation is 100 mm inward from the side edge. Three fields of view were photographed for each sample. The grain size of the largest grain in these photographs is regarded as the maximum grain size of Al inclusions (in terms of equivalent circle diameter). Macrosegregation and microsegregation were evaluated in the following manner.

(Macrosegregation)

The rolled slab is cut at arbitrary two positions such that each cut represents the upper and lower cross sections of the ingot. Two samples are taken from each of the cut slabs, representing the upper, middle, and lower parts of the ingot. The sampling position is surface and middle of thickness. The concentration of aluminum in each sample is determined. The maximum value of deviation from the aimed concentration (22%) is regarded as macrosegregation.

(Microsegregation)

Either of the two samples used for evaluation of macrosegregation was examined by an EPMA (electron probe microanalysis), with a beam (about 10 μm in diameter) scanning over an arbitrary length of 1 mm. Microsegregation was rated by judging whether or not fluctuation of aluminum concentrations is within 2%.

(Detection of Pores)

A reference sample was prepared as follows. A square bar, 50×50 mm, is cut out of the slab. It is reheated at 350° C. for 10 hours and then undergoes HIP, so that it is completely free of pores. A hole, 0.5 mm or 0.3 mm in diameter, is made by drilling at the center of the pore-free sample. The drilled reference sample is ultrasonically tested to determine the noise level that detects the 0.5-mm hole or 0.3-mm hole.

Each sample (20-mm thick slab) underwent ultrasonic test, and any sample that gave a noise higher than the predetermined one was regarded as having pores larger than 0.5 mm or 0.3 mm in diameter. The results of these tests are shown in Table 2 below, in which ⊚ denotes those samples containing pores smaller than 0.3 mm and ◯ denotes those samples containing only pores smaller than 0.5 mm.

TABLE 2 Dynamic Static Maximum Metallographic characteristics characteristics Macro- Macro- diameter of structure Tensile Tensile Sample segregation segregation Al inclusions α phase β phase strength Elongation strength Elongation No. (mass %) (mass %) (μm) (μm) (μm) Pores (MPa) (%) (MPa) (%) 1 3.5 3.1 5 8.9 0.011 ◯ 132.2 145.2 293 41.3 2 3.3 2.9 4 10 0.019 X 135.1 149.4 287.5 45.3 3 2.5 1.8 7 3.1 0.038 X 128.3 181.5 283.4 89.3 4 2.5 2.7 5 6.1 0.06 X 129.1 121.4 300.3 39.1 5 2.5 1.8 7 4.4 2.1 ◯ 123.5 203.1 289.3 83.1 6 2.5 1.8 7 6.1 5.2 ◯ 129.2 204.3 289.1 48.1 7 2.5 1.8 7 4.4 2.2 ◯ 130.5 204.1 289.9 87.4 8 2.5 1.8 7 4.2 1.6 ◯ 132.1 203.8 291.3 86.9 9 2.5 1.8 7 5.9 4.1 ◯ 127.5 134.2 294.3 42.5 10 3.4 2.8 7 4.3 1.3 ◯ 128.2 220.3 291.3 38.2 11 2.6 1.8 58 3.9 1.5 ◯ 126.4 201.2 299 45.3 12 1.1 1.2 4 3.8 0.8 ◯ 113.1 250.4 298.1 120.5 13 2.5 2.7 5 7.3 4.3 ◯ 127.5 121.4 300.3 38.6 14 2.5 1.7 30 4.3 1.5 ◯ 127.9 173.1 298.4 69.5 15 2.5 1.8 6 5.4 1.4 ◯ 128.4 119.3 299.3 33.6 16 2.5 1.5 60 4.2 1.3 ◯ 128.5 111.5 289.7 34.6 17 2.5 1.5 7 4.8 3.1 ⊚ 127.0 201.0 298.8 72.0

Tables 1 and 2 indicate the following. Samples Nos. 5, 7, 8, 12, and 14, which meet the requirements of the present invention, are excellent in both dynamic and static characteristics. Samples Nos. 5 and 14 are particularly excellent because they were prepared by using a water-cooled copper mold for the prescribed cooling rate to reduce macrosegregation and microsegregation. They exhibit good deforming characteristics even though they are made from a large ingot (180 kg).

However, samples Nos. 7 and 8, which meet the requirements of the present invention, are liable to forge cracking because of the low blooming (forging) temperature. This suggests that blooming (forging) should be carried out at 200° C. or above.

Sample No. 17 exhibits good deforming characteristics because it contains the α phase which has a grain size smaller than 5 μm and has no pores larger than 0.3 mm although it contains the β phase which has a grain size larger than 3 μm.

By contrast, samples Nos. 1, 2, 4, 6, 9, 10, 11, 13, 15, and 16, which do not meet any one of the requirements of the present invention, are poor in either static characteristics or dynamic characteristics, or are liable to forge cracking.

To be more specific, samples Nos. 1, 2, and 10 suffer remarkable macrosegregation and microsegregation on account of the low average cooling rates (1 and 2) and are poor in both static characteristics and dynamic characteristics. Sample No. 6 has a coarse metallographic structure and a low elongation on account of the excessively high temperature of warm working (rolling).

Sample No. 9 has a coarse metallographic structure and is poor in both static characteristics and dynamic characteristics on account of the excessively high temperature of blooming (forging).

Samples Nos. 11 and 16 have large Al inclusions and are poor in dynamic characteristics because casting was performed without the melt being isolated from the ambient atmosphere.

Sample No. 4 suffers noticeable microsegregation and has a coarse structure and is poor in both static characteristics and dynamic characteristics on account of the low reheating temperature and the omission of blooming (forging). Sample No. 13 is poor in both static characteristics and dynamic characteristics, with remarkable microsegregation and coarse metallographic structure, on account of the low reheating temperature. Sample No. 15 is poor in both static characteristics and dynamic characteristics on account of the large grain size of α phase.

Sample No. 3, which is a reference example in which the grain size of β phase is held below 0.05 μm as the present inventors proposed previously, is superior in both static characteristics and dynamic characteristics despite the presence of pores. 

1. A Zn—Al alloy excellent in rapid deformation properties containing 30 to 99% Zn, with the remainder being Al and inevitable impurities, which has a metallographic structure in which the α phase or α′ phase contains the β phase dispersed therein, the α phase or α′ phase has an average grain size no larger than 5 μm, the Al inclusions have a maximum equivalent circle diameter no larger than 50 μm and are free of pores no smaller than 0.5 mm in terms of equivalent circle diameter, and the macrosegregation of Al is less than 3.0% and the microsegregation of Al is less than 2.0%. (% means mass %.)
 2. The Zn—Al alloy excellent in rapid deformation properties as defined in claim 1, wherein the β phase has an average grain size no larger than 3 μm, preferably no larger than 0.1 μm.
 3. A method for producing a Zn—Al alloy excellent in rapid deformation properties, which comprises: a step of casting a molten Zn—Al alloy into a mold while isolating the melt from the ambient atmosphere, a step of cooling the mold after casting at an average cooling rate no lower than 0.25° C./s in the temperature range from 425 to 375° C. and no lower than 0.020° C./s in the temperature range from 275 to 250° C., a step of reheating which includes keeping hot at 350° C. or above and subsequent quenching, a step of blooming at 275° C. or below, and a step of warm working at 275° C. or below. 